The present invention relates generally to articles made with superalloys for use at high temperature and high stress. More particularly it relates to components of jet engines and turbines which are formed of superalloy materials and which are used in service at high temperatures and high applied stress.
It is known that superalloys including nickel-base and iron-base superalloys have been employed extensively in applications which require high strength at high temperature. The design of jet engines has in large part been determined by the properties which superalloys used as fabricating materials for components of the engine can display. As the properties of the alloys are improved the design of the jet engine improves and greater thrust to weight ratios are achieved. Generally higher temperature operation results in greater fuel efficiency for such engines and the drive for higher operating temperatures and for superalloy materials which can operate at such higher temperatures is a continuous design criteria in fabrication of more and more efficient jet engines. The need for higher temperature capability in high strength superalloys continues as efforts are made to continue to improve operating performance for jet engines.
Many metallurgical advances have assisted in improving high strength superalloys. These have included the increase in the precipitate volume fraction for the gamma precipitate strengthening agent of such alloys. Also improvements have been made through powder metallurgy and through the use of isothermal forging. Improvements in the alloy temperature capability of superalloys have been achieved.
It has also been recognized that not all components of a jet engine are subject to the same operating conditions and that different metallurgical compositions may be employed in different components of the engine to best suit the needs for that component. Also it has been found that different heat treatments of a single composition can result in different combinations of properties and such different heat treatments have been employed depending on the use and application of the engine part and the function which it serves in the overall engine.
However, there are some parts where tradeoffs have been made in properties because the part is large enough so that the engine operating conditions over the full extent of the part are not uniform. In other words, certain large pieces which are installed in an engine encounter different temperatures and different property requirements in service from one portion of the component to another. As is known, conventional thermal processing imparts an identical thermal history to each portion of a component and generates the same microstructure and properties throughout the whole component. Accordingly, for such large components it is necessary to sacrifice a property at one location of the component in order to obtain an acceptable property at another location.
For example, among the many critical mechanical properties required by highly stressed components, crack growth resistance and high temperature rupture life are highly desired properties. Such properties are needed, for example, in engine disks which rotate at high speeds and result in the application of high stress to portions of the disk and particularly to the outer portions of the disk. A number of improvements have been made recently in crack growth resistance of superalloys employed in forming such disks and also in the enhancement of high temperature rupture life for such disks.
The prior art heat treatment practice for high strength superalloys employs subsolvus annealing and rapid cooling in order to generate a fine grain structure and high strength. Subsolvus annealing is an annealing at a temperature below the solvus temperature, i.e., the temperature at which all .gamma.' strengthening precipitate goes into solution. The subsolvus annealing is also known as partial solutioning. The conventional cooling rate after annealing is the fastest rate which is possible providing that no quench defect is induced in the part. Alloys receiving this kind of conventional heat treatment show good tensile and fatigue strength but their high temperature rupture life is comparatively short. Improvements in high temperature rupture life can be accomplished by supersolvus annealing to generate a large grain size.
An illustration of the different results and different properties which are imparted to a single base superalloy is illustrated in the data which is plotted in FIG. 1.
Referring now first to FIG. 1, it is evident from the plotted data that for a sample of Rene 95 prepared by powder metallurgy techniques a 100 hour rupture life test was performed. The test applied a level of stress to the sample as plotted in the ordinate and the sample was heated to the various temperatures noted in the abscissa of the plot of temperature in degrees Fahrenheit. At 1000.degree. F. the material which had been annealed at subsolvus temperature was stressed at about 190 ksi and can endure for 100 hours with this stress at that temperature. At 1100.degree. F. the stress which can be endured for 100 hours is about 170 ksi. At 1400.degree. F. the stress which can be endured by the subsolvus annealed material is about 50 ksi.
In contrast with the results obtained for the subsolvus annealed material, the supersolvus annealed material can be seen to have a much higher stress at 1400.degree. and, in fact, about 90 ksi. The rupture strength for 100 hour life is plotted as a function of temperature for subsolvus and supersolvus annealed materials, respectively. The two rupture curves show a crossover at about 1240.degree. F. At a stress of about 80 ksi the supersolvus annealing offers about +100.degree. F. temperature capability over the subsolvus annealed material. In contrast the latter, that is the subsolvus annealed material, has a much higher strength than the supersolvus annealed material at low temperatures.
The cooling rate after annealing also affects alloy properties significantly. Generally, higher cooling rates result in higher tensile strengths for the same composition. FIG. 2 is a plot of the strength of a Rene 95 sample prepared by powder metallurgy plotted as the ordinate against the cooling rate employed in cooling the sample as the abscissa. The yield and tensile strengths were measured on samples at 1200.degree. F. and results are plotted in FIG. 2. It is noted that as the cooling rate in degrees Centigrade per minute increases, that both the yield strength and the tensile strength of the samples increases based on the different cooling rates after annealing and based particularly on the increased cooling rate. In other words, it has been found that the higher the cooling rate employed in affecting the cooling of the sample, the higher the strengths that are obtained. This same finding of increased strength with increased cooling rate is found to be the case for both subsolvus annealed and supersolvus annealed samples.
The most striking effect of cooling rate on the properties of a supersolvus annealed sample are identified to be the fatigue crack growth resistance under time dependent conditions. This is described in the copending application Ser. No. 907,550 filed Sept. 15, 1986. An example to illustrate the time dependence of fatigue crack propagation is given in FIG. 3. In FIG. 3 the rate of crack propagation, da/dN, in inches per cycle is plotted against the cooling rate in .degree. C. per minute. The fatigue crack growth rate for the samples was measured at 1200.degree. F. by employing three cycle waveforms as follows: The first was a 3 second sinusoidal cycle. The second was a 180 second sinusoidal cycle. The third was a 177 second hold at the maximum load of a 3 second sinusoidal cycle. In these tests the maximum to minimum load ratio was set at R=0.05. From the data plotted in FIG. 3, it is evident that for a given cyclic stress intensity, .DELTA.K=30 ksi.sqroot.in, the crack growth da/dN, of fast cooled Rene 95 increases dramatically as the cycle frequency decreases or as the hold time is applied. From the graph of FIG. 3, the lowest increase in crack growth rate is seen to be for the 3 second cycle. The 180 second cycle is seen to be dramatically higher than that of the 3 second cycle and the 3 second cycle with 170 second hold at maximum stress is seen to have a dramatically greater crack growth propagation rate. For example, and with reference again to FIG. 3, at a cooling rate of 750.degree. C. per minute the da/dN for the 180 second cycle is about 0.006 inches per cycle while that for the 3 second cycle is about 0.00004 inches per cycle.
The time dependence of the da/dN, that is of the crack growth rate, becomes less and less when the cooling rate after supersolvus annealing is reduced to greater and greater extents. These, and again from the data plotted in FIG. 3, at a cooling rate of 100.degree. C. per minute the da/dN for the 180 second cycle is about 0.00002 inches per cycle while that for the 3 second cycle is about 0.00004 inches per cycle.
From the foregoing it is evident that the optimum thermal processing for high temperature capability with resistance to crack propagation or growth is completely different from that for high strength.